JP4858004B2 - High-strength steel sheet with excellent ductility and deep drawability and method for producing the same - Google Patents

Description

  This invention is useful for steel sheets for automobiles, etc., and has a high tensile strength TS of 440 MPa or higher, an average r value (Rankford value) of 1.2 or higher, excellent deep drawability, and excellent ductility, that is, tensile strength TS The strength-ductility balance represented by the product TS x El of the total elongation El and the strength-ductility balance and the strength represented by the product TS x U-El of TS and uniform elongation U-El- It relates to the manufacturing method.

In recent years, in order to regulate CO ₂ emissions from the viewpoint of global environmental conservation, improvement in fuel efficiency of automobiles has been demanded. In addition, in order to ensure the safety of passengers in the event of a collision, safety improvements centering on the collision characteristics of automobile bodies are also required. For this reason, the weight reduction and reinforcement of the automobile body are being actively promoted.

  In order to satisfy the weight reduction and strengthening of automobile bodies at the same time, it is said that it is effective to increase the strength of component materials and reduce the plate thickness within a range where rigidity is not a problem. Are actively used in automotive parts. In particular, since the weight reduction effect increases as the strength of the steel sheet used increases, in the automobile industry, for example, steel sheets having a TS of 440 MPa or more are used as inner and outer panel material.

  On the other hand, since many automotive parts made of steel plates are manufactured by press forming, the steel plates for automobiles require excellent press formability. In general, high-strength steel sheets are significantly inferior in formability, particularly deep drawability and ductility, compared to ordinary mild steel sheets, so in promoting weight reduction of automobile bodies, TS is 440 MPa or more, preferably 500 MPa or more, More preferably, a high-strength steel sheet is required that is 590 MPa or more, an average r value that is an index of deep drawability is 1.2 or more, and TS × El is 20000 MPa ·% or more. Furthermore, TS × U-El is desired to be 10000 MPa ·% or more for parts where overhang is important.

  As a means to increase the strength while having such a high r value, based on IF (Interstitial Free) steel, in which Ti or Nb is added to extremely low carbon steel and solid solution carbon or solid solution nitrogen is fixed, it is based on this. There is a method of adding solid solution strengthening elements such as Si, Mn, and P. For example, Patent Document 1 contains C: 0.002 to 0.015%, Si: 1.2% or less, Mn: 0.04 to 0.8%, P: 0.03 to 0.10%, and Nb C × 3 to (C × 8 + 0.020 )% (Where C represents the content of element C), TS is 340 to 460 MPa, average r value is 1.7 or more, El is 36% or more, and non-aging A high-tensile cold-rolled steel sheet having excellent properties is disclosed. However, it is difficult to stably produce a steel plate with a TS of 440 MPa or more using such an ultra-low carbon steel as a raw material, and when trying to produce a steel plate with a TS of 500 MPa or more, the amount of alloying elements increases, Problems such as deterioration of surface appearance, deterioration of plating properties, and manifestation of secondary processing brittleness occur. Furthermore, since the r value deteriorates when a large amount of solid solution strengthening element is added, the r value decreases as the strength increases.

  In order to increase the strength of the steel sheet, there is a method by strengthening the structure in addition to the method by solid solution strengthening as described above. For example, there is a DP (Dual-Phase) steel sheet having a composite structure composed of a soft ferrite phase and a hard martensite phase. DP steel is generally characterized by excellent ductility and excellent TS x El, and low yield ratio, that is, low yield stress for tensile strength and excellent shape freezing properties during press forming. However, the r value is low and the deep drawability is poor. This is because there is a martensite phase that does not contribute to the r value in terms of crystal orientation, or solid solution C, which is essential for the formation of the martensite phase, inhibits the formation of {111} recrystallized texture effective for increasing the r value It is said to do.

As a technique for improving the r value of DP steel sheet, Patent Document 2 includes C: 0.05 to 0.15%, Si: 1.50% or less, Mn: 0.30 to 1.50%, P: 0.030% or less, S: 0.030% or less, sol .Al: 0.020 to 0.070%, N: 0.0020 to 0.0080%, steel consisting of the remaining Fe and inevitable impurities is hot rolled at the finish rolling exit temperature above the Ar ₃ transformation point at a coiling temperature of 600 ° C or less, After cold rolling at a rolling reduction of 40% or more, batch annealing (box annealing) is performed, and then heated to 700 to 800 ° C in a continuous annealing furnace to form a composite structure, and after water quenching, at 200 to 500 ° C A method for producing a high-strength cold-rolled steel sheet excellent in deep drawability for tempering is disclosed.

  In Patent Document 3, C: 0.20% or less, Si: 1.0% or less, Mn: 0.8 to 2.5%, Sol.Al: 0.01 to 0.20%, N: 0.0015 to 0.0150%, P: 0.10% or less, the balance is substantially In particular, steel made of Fe is hot-rolled, cold-rolled, then box-annealed in a temperature range of 650 to 800 ° C, and then heated to 600 to 850 ° C in a continuous annealing furnace. A method for producing a high-tensile cold-rolled steel sheet excellent in deep drawability and shape properties that is cooled at a cooling rate of s is disclosed. In this method, Mn is concentrated from the ferrite (α) phase to the austenite (γ) phase during soaking of the box annealing, and this Mn enriched phase is preferentially used as the γ phase during the subsequent soaking of the continuous annealing. Sometimes complex tissues are formed.

Patent Document 4 contains C: 0.003 to 0.03%, Si: 0.2 to 1%, Mn: 0.3 to 1.5%, Al: 0.01 to 0.07%, Ti: 0.02 to 0.2%, and an atomic concentration ratio (effective Ti ) / (C + N) (where Ti, C, and N represent the content of each element) were hot-rolled and cold-rolled steel, and after cold rolling, Ac _{1 was obtained} by continuous annealing. Disclosed is a method for producing a composite structure type high-tensile cold-rolled steel sheet that is excellent in both deep drawability and shape by rapid cooling at an average cooling rate of 30 ° C / s or higher after heating to a temperature range of not less than the transformation point and not more than 900 ° C. Yes. Actually, the steel with the composition of 0.012% C-0.32% Si-0.53% Mn-0.03% P-0.051% Ti in mass% was cold-rolled and then heated to 870 ° C which is the α-γ2 phase region, and then in the jet water That is, by cooling at an average cooling rate of 100 ° C./s, a composite-structure high-tensile cold-rolled steel sheet having an average r value of 1.6, TS of 492 MPa, and El of 36% is obtained.

  In Patent Document 5, C: 0.01 to 0.080%, Si: 2.0% or less, Mn: 3.0% or less, P: 0.10% or less, S: 0.02% or less, Al: 0.005-0.20%, N: 0.02% or less and V: 0.01 to 0.5%, and V and C satisfy the relationship of 0.5 × C / 12 ≦ V / 51 ≦ 3 × C / 12 (where V and C represent the content of each element) Also disclosed is a high-strength cold-rolled steel sheet having excellent structure and deep drawability, the balance being substantially composed of Fe and inevitable impurities. In this steel sheet, before recrystallization annealing, C in the steel is precipitated as V-based carbides to reduce the solute C as much as possible to increase the r value, and then continuously heat to the α-γ2 phase region by continuous annealing. V-type carbides are dissolved by C to concentrate C in the γ phase, and a martensite phase is formed in the subsequent cooling process.

  In Patent Document 6, C: 0.03-0.25%, Si: 0.001-3.0%, Mn: 0.01-3.0%, P: 0.001-0.06%, S: 0.05% or less, Al: 0.005-0.3%, N: 0.001 Contain ~ 0.030%, balance is Fe and inevitable impurities, average r value is 1.3 or more, and contains 3 ~ 100% of at least one of bainite phase, martensite phase and austenite phase in the structure A high-strength steel sheet having excellent deep drawability is disclosed. This steel sheet is annealed at a heating rate of 4 to 200 ° C./h in order to achieve a high r value, that is, box annealing, and then subjected to a short heat treatment to form a bainite phase, a martensite phase, an austenite phase, etc. Applied and manufactured.

In Patent Document 7, C: 0.010 to 0.050%, Si: 1% or less, Mn: 1.0 to 3.0%, P: 0.005 to 0.1%, S: 0.01% or less, Al: 0.005 to 0.5%, N: 0.01% Hereinafter, Nb: 0.01 to 0.3%, and the content of Nb and C is (Nb / 93) / (C / 12) = 0.2 to 0.7 (where Nb and C represent the content of each element) ), The balance being substantially composed of Fe and inevitable impurities, and having a composite structure including a ferrite phase with an area ratio of 50% or more and a martensite phase with an area ratio of 1% or more. A high-strength steel sheet excellent in deep drawability having an average r value of 1.2 or more is disclosed. This steel plate is hot rolled from a slab having the above composition at a finish rolling exit temperature of 800 ° C or higher, wound at a winding temperature of 420 to 720 ° C, and after cold rolling, annealed at an annealing temperature of 800 to 950 ° C. And the temperature range from the annealing temperature to 500 ° C. is cooled at an average cooling rate of 5 ° C./s or more.
JP-A-56-139654 Japanese Patent Publication No.55-10650 JP 55-100934 Japanese Patent Publication No. 1-35900 Japanese Patent Laid-Open No. 2002-226941 JP 2003-64444 A JP 2005-120467 A

  However, these conventional techniques have the following problems.

  Patent Documents 2, 3, and 6: Box annealing is required, and productivity is remarkably inferior, and there are many manufacturing problems such as adhesion of steel plates, generation of temper collars, and reduction in the life of the furnace body inner cover.

Patent Documents 2 and 4: Special equipment such as water quenching and jet water cooling is required, resulting in an increase in cost and a problem of surface treatment properties of the steel sheet.
Patent Document 5: Ductility is inferior due to the influence of VC and the like, and excellent TS × El and TS × U-El cannot be secured.

  Patent Document 7: Depending on manufacturing conditions, it may not be possible to secure excellent TS × El and TS × U-El.

  The present invention does not require low-productivity box annealing or special equipment, and TS of 440 MPa or more, average r value of 1.2 or more, TS × El of 20000 MPa ·% or more, TS × U of 10000 MPa ·% or more It is an object of the present invention to provide a high-strength steel plate in which -El is stably obtained and a method for producing the same.

  As a result of diligent studies to solve the above-mentioned problems, the present inventors set the C amount to 0.020 to 0.050%, regulates the Nb amount according to this C amount, and in the matrix composed of the ferrite phase. Furthermore, it has been found that it is effective to uniformly disperse the second phase containing the martensite phase at an appropriate area ratio.

The present invention was made based on such findings, and in mass%, C: 0.020 to 0.050%, Si: 1.0% or less, Mn: 1.0 to 2.5%, P: 0.005 to 0.1%, S: 0.01% Hereinafter, Al: 0.005 to 0.5%, N: 0.01% or less, Nb: 0.010 to 0.3%, Nb and C content satisfy the following formula (1), the balance is Fe and inevitable impurities Nmax defined by the following: a second phase containing a martensite phase with an area ratio of 1% or more with respect to the entire structure is contained in a matrix composed of a ferrite phase and having a composition of The present invention provides a high-strength steel sheet excellent in ductility and deep drawability, characterized in that it has a structure in which is 5 or more and an average r value is 1.2 or more.
(Nb / 93) / (C / 12) = 0.2 to 0.7 (1)
In the formula (1), each element symbol represents the content (% by mass) of each element,
Nmax is a distribution diagram of the center of gravity of the second phase obtained by image processing the image observed by the scanning electron microscope of the steel plate thickness cross section, and draws a circle with the ferrite average grain diameter as the radius around an arbitrary center of gravity, This is the number of centroids that appears most frequently when the operation for obtaining the number of centroids included in a circle other than the center of gravity of the circle is performed on all the centroids recognized in the observed image.

  Further, in addition to the above composition, Cr: 0.5% or less can also be contained by mass%.

  The high-strength steel sheet of the present invention is in mass%, C: 0.020 to 0.050%, Si: 1.0% or less, Mn: 1.0 to 2.5%, P: 0.005 to 0.1%, S: 0.01% or less, Al: 0.005 to 0.5 %, N: 0.01% or less, Nb: 0.010 to 0.3%, and the content of Nb and C satisfies the above formula (1), and the balance is composed of Fe and inevitable impurities. Hot rolled at a finish rolling exit temperature of 800 ° C. or higher, wound into a hot rolled steel sheet at a winding temperature of 400 to 720 ° C., then cold rolled at a rolling reduction of 40% or higher to obtain a cold rolled steel sheet, The cold-rolled steel sheet is heated to an annealing temperature of 800 to 950 ° C., then the temperature range from the annealing temperature to 580 ° C. is cooled at an average cooling rate of 5 to 30 ° C./s, and an overaging temperature of 480 to 580 ° C. After holding for 30 to 600 s, it can be produced by a method of cooling at least up to 100 ° C. at an average cooling rate of 5 ° C./s or more.

In the method of the present invention, in addition to the above composition, the it is et al., In mass%, Cr: may contain 0.5% or less.

  According to the present invention, TS ≧ 440 MPa, high strength, excellent average r value of 1.2 or more and deep drawability without requiring low productivity box annealing and special equipment, TS × 20,000 MPa ·% or more El, TS × U-El of 10000 MPa ·% or more and high-strength steel sheets with excellent ductility can be manufactured stably. By applying the high-strength steel sheet of the present invention to automobile parts, it is possible to increase the strength of parts that have been difficult to press-form so far, and it is possible to sufficiently achieve collision safety and weight reduction of automobile bodies. Moreover, the high-strength steel sheet of the present invention is applicable not only to automobile parts but also to home appliance parts and pipe materials.

  Details of the present invention will be described below. The “%” below represents “% by mass” unless otherwise specified.

1) ingredients
C: 0.020 ~ 0.050%
C is effective for increasing the strength and is an important element in the present invention together with Nb described later. In order to obtain TS of 440 MPa or higher and high El and U-El, as described later, it is necessary to uniformly disperse the second phase including the martensite phase in the matrix of the ferrite phase. In order to increase the strength by forming a martensite phase, the C content needs to be 0.020% or more. In particular, in order to obtain a TS of 590 MPa or more, it is desirable that the C content be 0.025% or more. On the other hand, excessive addition hinders the development of {111} recrystallization texture during annealing, and an average r value of 1.2 or more cannot be obtained, so the upper limit of C content is 0.050%.

Si: 1.0% or less
In addition to the effect of solid solution strengthening, Si promotes α transformation and increases the C content in the untransformed γ phase, making it easier to form a composite structure of the ferrite phase and the second phase including the martensite phase. Has an effect. In order to obtain the above effect, the Si content is preferably 0.01% or more, and more preferably 0.05% or more. On the other hand, if the Si content exceeds 1.0%, a surface defect called red scale occurs during hot rolling, which deteriorates the surface appearance of the steel sheet. Cause uneven plating. Therefore, the Si content is 1.0% or less, preferably 0.7% or less.

Mn: 1.0-2.5%
Mn is effective for increasing the strength and has the effect of lowering the critical cooling rate at which the second phase including the martensite phase is obtained, and promotes the formation of the second phase during the cooling of the annealing. Therefore, the amount is preferably adjusted according to the required strength level and the cooling rate of annealing. Mn is an element effective for preventing hot cracking due to S. From such a viewpoint, the amount of Mn needs to be 1.0% or more, preferably 1.2% or more. On the other hand, if the Mn content exceeds 2.5%, the r value and weldability deteriorate, so the upper limit of the Mn content is set to 2.5%.

P: 0.005-0.1%
P has a solid solution strengthening effect. However, if the amount of P is less than 0.005%, not only the effect does not appear, but also the dephosphorization cost at the time of steelmaking increases. Therefore, the P content is 0.005% or more, preferably 0.01% or more. On the other hand, if the amount of P exceeds 0.1%, P segregates at the grain boundaries and deteriorates the secondary work brittleness resistance and weldability. In addition, during alloying after hot dip galvanizing, P suppresses the diffusion of Fe at the interface between the plating layer and the steel sheet and degrades the alloying processability. Therefore, alloying at high temperatures is required, and powdering and Plating peeling such as chipping is likely to occur. Therefore, the upper limit of the P content is 0.1%.

S: 0.01% or less
In addition to causing hot cracking, S is present as an inclusion in steel and deteriorates various properties of the steel sheet. Therefore, the S amount needs to be 0.01% or less, but is preferably reduced as much as possible.

Al: 0.005-0.5%
In addition to being useful as a solid solution strengthening element and deoxidizing element for steel, Al has the effect of improving the normal temperature aging resistance by precipitating solid solution N as AlN. Al is also useful for adjusting the temperature in the α-γ2 phase region as a ferrite-forming element. Therefore, the Al amount needs to be 0.005% or more. On the other hand, if the Al content exceeds 0.5%, alloy costs increase and surface defects are induced, so the upper limit of Al content is 0.5% or less, preferably 0.1% or less.

N: 0.01% or less
When N is present in a large amount, the room temperature aging resistance is deteriorated, so that a large amount of Al or Ti needs to be added accordingly. Therefore, the upper limit of the N amount needs to be 0.01%, but it is preferable to reduce it as much as possible.

Nb: 0.010-0.3%
Nb is the most important element in the present invention, and contributes to increasing the r value by refining the structure of the hot-rolled steel sheet, or by precipitating it as NbC in the hot-rolled steel sheet and reducing the amount of dissolved C. From such a viewpoint, the Nb amount needs to be 0.010% or more. On the other hand, in order to form the second phase including the martensite phase during the cooling process of annealing, solid solution C is required, and for that purpose, the Nb content needs to be 0.3% or less.

(Nb / 93) / (C / 12): 0.2 to 0.7 (where Nb and C represent the content of each element)
As described above, it is necessary to control (Nb / 93) / (C / 12) representing the atomic concentration ratio of Nb and C as follows after controlling the amounts of C and Nb. When (Nb / 93) / (C / 12) is less than 0.2, the effect of refining the structure of the hot-rolled steel sheet is small, and the amount of solute C increases, which is advantageous for increasing the r value. Inhibits formation. On the other hand, if (Nb / 93) / (C / 12) exceeds 0.7, the amount of C necessary to form the second phase including the martensite phase is prevented from being present in the steel, and finally the martensite. The second phase including the site phase is not formed. Therefore, in order to achieve both high r value and high ductility, (Nb / 93) / (C / 12) needs to be 0.2 to 0.7, preferably 0.2 to 0.5.

  The balance is Fe and inevitable impurities. Here, the inevitable impurities include 0.01% or less of Sb, 0.1% or less of Sn, 0.01% or less of Zn, 0.1% or less of Co, and 0.05% or less of Mo. In particular, Mo makes the dispersion state of the second phase including the martensite phase non-uniform and excellent uniform elongation cannot be obtained. Therefore, the amount of Mo as an impurity needs to be limited to 0.05% or less.

  In order to achieve the object of the present invention, the above components are sufficient. However, Ti: 0.1% or less and Cr: 0.5% or less can be further contained for the following reasons.

Ti: 0.1% or less
Ti has the same affinity as Al or higher affinity for N than Al, and has the effect of precipitating solute N. In order to obtain this effect, the Ti content is preferably 0.005% or more. However, if it exceeds 0.1%, not only will the cost increase, but the amount of solute C required to form the second phase including the martensite phase will be reduced by the formation of TiC. Therefore, the Ti content is 0.1% or less. Further, Ti preferentially bonds with S and N in the steel, and then bonds with C. Considering the decrease in Ti yield due to the formation of inclusions in steel, etc., when Ti content exceeds (Ti / 48) / {(S / 32) + (N / 14)} exceeds 2.0, S and N The effect of Ti addition of fixing is saturated, and on the contrary, the effect of promoting the formation of TiC and preventing the solute C from remaining in the steel is increased. Therefore, the Ti content should satisfy (Ti / 48) / {(S / 32) + (N / 14)} ≦ 2.0 in relation to the contents of S and N that are preferentially bonded in steel. Is preferred. Here, Ti, S, and N in the relational expression are the contents (mass%) of each element.

Cr: 0.5% or less
Cr, like Mn, is effective for increasing the strength and has the effect of lowering the critical cooling rate at which a second phase including a martensite phase is obtained, and promotes the formation of the second phase during annealing cooling. . In order to obtain this effect, the Cr content is preferably 0.05% or more. On the other hand, if the Cr content exceeds 0.5%, the r value is lowered, so the Cr content is 0.5% or less.

  Furthermore, B, Ca, REM, and the like can be contained as long as they are within the normal steel composition range. For example, B that improves the hardenability of steel is contained in the range of 0.003% or less, and at least one element of Ca and REM effective for the shape control of sulfide inclusions is contained in the range of 0.01% or less. it can.

2) Organization
2-1) Area ratio of the second phase including martensite phase of 1% or more in terms of the area ratio relative to the entire structure: 15% or less
In order to obtain a high-strength steel sheet having a TS of 440 MPa or more, good TS × El, TS × U-El and deep drawability, a martens of 1% or more in the area ratio of the entire structure in the matrix composed of ferrite phase The second phase including the site phase needs to have a structure including 15% or less in area ratio. Here, the entire structure means the entire steel sheet structure including the matrix composed of the ferrite phase and the second phase. In the present invention, an average r value of 1.2 or more can be obtained by using a ferrite phase in which a {111} recrystallized texture is developed as the matrix (structure main body). Here, the ferrite phase means a polygonal ferrite phase or a bainitic ferrite phase having a high dislocation density transformed from a γ phase.

  In order to reliably obtain a TS of 440 MPa or more, the area ratio of the martensite phase needs to be 1% or more in terms of the area ratio with respect to the entire structure, and more preferably 3% or more. As described above, the second phase including the martensite phase has an average r value of 1.2 or more by using a ferrite phase in which a {111} recrystallized texture is developed as a matrix. Must be less than%. If the area ratio of the second phase exceeds 15%, that is, the area ratio of the ferrite phase is less than 85%, it is difficult to ensure an average r value of 1.2 or more. The structure of the second phase need not be specified except for the martensite phase described above, and may be a structure containing a bainite phase, a cementite phase, or the like. Here, the area ratio of the ferrite phase and the above-mentioned second phase and the average ferrite grain size described later were observed by SEM for the plate thickness section parallel to the rolling direction, and image analysis software (Media Cybernetics “Image-ProPLUSVer. 4.0.0.11).

2-2) Dispersibility of the second phase: Nmax ≧ 5
In order to obtain TS × El of 20000 MPa ·% or more and TS × U-El of 10000 MPa ·% or more stably, it is necessary to further increase the Nmax as defined above to 5 or more in addition to the above-mentioned organizational regulations. is there.

  Specifically, Nmax is measured as follows. First, the SEM observation image of the plate thickness cross section parallel to the rolling direction as shown in FIG. 1 (a), by image processing, the ferrite phase as shown in (b) and the second phase as shown in (c) The center of gravity coordinates where the first moment around the second phase is 0 are obtained using the above image analysis software, and the distribution map of the center of gravity of the second phase as shown in (d) is obtained. In FIG. 1, examples of steel plates Nos. 1, 2 and 16 of examples described later are shown, but it can be seen that the distribution of the center of gravity of the second phase differs greatly depending on the steel plates. Next, select an arbitrary center of gravity from the distribution diagram of the center of gravity of the second phase shown in (d) of FIG. 1, and as shown in FIG. 2, the ferrite obtained using (b) of FIG. A circle with the average particle diameter D as the radius is drawn, and the number of centroids contained in the circle other than the centroid at the center of the circle (6 in the case of FIG. 2) is obtained. At this time, regarding the second phase existing at the edge of the observed image, a periodic boundary condition is assumed and the number of centroids is obtained as follows. That is, when the center of gravity that is the center of the circle is present at the edge, the same field of view is arranged so as to surround the observation image being measured, and the number of centers of gravity contained in the circle having the average ferrite particle diameter D is obtained. Note that the centroid existing on the circle having the average ferrite particle diameter D is not counted as the number of centroids. This operation is repeated with the other centroid as the center of the circle, and the number of centroids that appears most frequently is Nmax. Fig. 3 shows the relationship between the number of centroids and the frequency (ratio of the number of centroids where each center of gravity appears relative to the total number of measurements) in steel plates No. 1, No. 2 and No. 16. In .2, Nmax is 5 or more, and in steel plate No. 16, Nmax is 4. As shown in Table 3 to be described later, excellent TS × El and TS × U-El were obtained for steel plates No. 1 and No. 2. The inventors have conducted such studies on various steel plates and found that excellent TS × El and TS × U-El can be obtained by setting Nmax to 5 or more.

  The reason why excellent TS × El and TS × U-El can be obtained when Nmax is 5 or more is considered as follows. That is, one ferrite grain is completely contained in a circle whose radius is the average ferrite grain diameter D. Assuming that the grain boundary triple point is the optimal location of the second phase, when Nmax is 5 or more, the second phase is evenly distributed at the triple point, and the second phase is uniformly dispersed in the ferrite phase matrix. Therefore, U-El is improved, and excellent TS × El and TS × U-El are obtained.

The high-strength steel sheet of the present invention satisfies the above components and steel microstructure, and has an average r value of 1.2 or more. The steel sheet of the present invention having such an average r value is determined by X-ray diffraction at the 1/4 thickness position of the steel sheet, and the (222) plane, (200) plane, (110) plane parallel to the plane And (310) plane X-ray diffraction integrated intensity ratios P ₍₂₂₂₎ , P ₍₂₀₀₎ , P ₍₁₁₀₎ , P ₍₃₁₀₎ are P ₍₂₂₂₎ / {P ₍₂₀₀₎ + P ₍₁₁₀₎ + P ₍₃₁₀₎ } ≧ 1.5 is preferably satisfied, and more preferably P ₍₂₂₂₎ / {P ₍₂₀₀₎ + P ₍₁₁₀₎ + P ₍₃₁₀₎ } ≧ 2.0.

  Conventionally, the r value is high when the {111} surface has a texture parallel to the plate surface, but the r value is low when the {110} surface or {100} surface is parallel to the plate surface. It has been. Here, the details are not yet clear, but the {310} plane is a texture that lowers the r-value, like the {100} and {110} planes, although the strength is low. We know that this also contributes to higher r-values. In addition, in the steel sheet of the present invention, because Nb is added, the rolling reduction in the non-recrystallized γ region at the time of hot rolling is high, the precipitation of fine NbC, and the presence of C not precipitated and fixed as NbC, etc. However, it is thought that it contributes to the reduction of {310} surface strength.

In addition, {111} texture means that the <111> direction of the crystal is oriented in the direction perpendicular to the steel plate surface. From the Bragg reflection condition, in the case of α-Fe with a body-centered cubic structure, diffraction on the {111} plane does not occur on the (111) plane, but occurs on the (222) plane. The value of (222) plane (P ₍₂₂₂₎ ) was used. Therefore, the high strength of the (222) plane corresponds to the development of the {111} texture. For the same reason, the value of (200) plane (P ₍₂₀₀₎ ) was used for {100} plane. Here, the X-ray diffraction integrated intensity ratio is a relative intensity based on the X-ray diffraction integrated intensity of a non-directional standard sample (irregular sample). The X-ray diffraction may be either an angle dispersion type or an energy dispersion type, and the X-ray source may be a characteristic X-ray or a white X-ray. The measurement surface is desirably measured from 7 to 10 surfaces (110) to (420) which are the main diffraction surfaces of α-Fe.

3) Manufacturing Method In the manufacturing method of the present invention, first, a steel slab having the above composition (hereinafter simply referred to as slab) is hot-rolled at a finish rolling exit temperature FT of 800 ° C. or higher, and rolled at 400 to 720 ° C. A hot rolled steel sheet is wound at a coiling temperature CT.

  At this time, the slab to be used is preferably manufactured by a continuous casting method to prevent macro segregation of components, but may be manufactured by an ingot-making method. The continuous casting method may be a thin slab casting method. Also, in order to hot-roll the slab, in addition to the conventional method in which the slab is once cooled to room temperature and then reheated and rolled, a method of hot rolling immediately after continuous casting, or a hot piece without cooling to room temperature. Energy-saving processes such as the method of charging in a heating furnace and rolling can be applied without problems.

  The slab heating temperature is preferably low because the precipitates are coarsened to develop a {111} recrystallized texture and improve deep drawability. However, if the heating temperature is less than 1000 ° C., the rolling load increases and the risk of trouble during hot rolling increases, so the slab heating temperature is preferably set to 1000 ° C. or higher. In order to prevent an increase in scale loss accompanying an increase in oxidized weight, the slab heating temperature is preferably 1300 ° C. or lower. The slab is subjected to hot rolling for rough rolling and finish rolling.

  Here, the slab is made into a sheet bar by rough rolling. The conditions for rough rolling are not particularly limited, and may be performed according to a conventional method. Further, when the heating temperature of the slab is lowered, it is preferable to heat the sheet bar using a sheet bar heater from the viewpoint of preventing troubles during rolling.

  The sheet bar is a hot-rolled sheet by finish rolling. At this time, in order to obtain a fine hot-rolled steel sheet structure so that excellent deep drawability can be obtained after cold rolling and recrystallization annealing, FT needs to be 800 ° C. or higher. If the FT is less than 800 ° C, the structure of the hot-rolled steel sheet has a processed structure, and the {111} recrystallized texture does not develop after cold rolling and annealing, and it is difficult to obtain a high r value. The load is also high. On the other hand, when the FT exceeds 980 ° C, the structure of the hot-rolled steel sheet becomes coarse, which may hinder the formation and development of {111} recrystallized texture after cold rolling and annealing, and may not obtain a high r value. FT is preferably 980 ° C. or lower.

  Moreover, in order to reduce the rolling load at the time of hot rolling or to make the shape and characteristics of the steel sheet uniform, lubrication rolling can be performed between some or all passes of finish rolling. The coefficient of friction during lubrication rolling is preferably in the range of 0.10 to 0.25. Furthermore, from the viewpoint of operational stability of hot rolling, it is preferable to apply a continuous rolling process in which sheet bars are joined and finish-rolled continuously.

  The hot-rolled steel sheet after hot rolling is wound, but the CT needs to be 400 to 720 ° C, preferably 550 to 680 ° C in order to refine the structure of the hot-rolled steel sheet and precipitate NbC. When CT exceeds 720 ° C, the crystal grains of the hot-rolled steel sheet become coarse, leading to a decrease in strength and a decrease in r value. Further, when the CT is less than 400 ° C., it is difficult for NbC to precipitate, and it becomes difficult to ensure deep drawability.

  Next, the hot rolled steel sheet is made into a cold rolled steel sheet by cold rolling. Note that the hot-rolled steel sheet is preferably pickled to remove scale before cold rolling. Pickling may be performed under normal conditions. If the rolling reduction of cold rolling is less than 40%, {111} recrystallized texture does not develop and it is difficult to obtain excellent deep drawability, so it is necessary to make it 40% or more, preferably 50% or more. There is. On the other hand, in the present invention, in the range up to 90%, the r value increases as the rolling reduction increases, but when the rolling reduction exceeds 90%, not only the effect is saturated, but also the load on the roll during rolling increases. Therefore, the upper limit of the rolling reduction is preferably 90%.

  The cold-rolled steel sheet is heated to an annealing temperature of 800 to 950 ° C., then cooled in the temperature range from the annealing temperature to 580 ° C. at an average cooling rate of 5 to 30 ° C./s, and with an overaging temperature of 480 to 580 ° C. After holding for 30 to 600 s, annealing is performed to cool to at least 100 ° C. at an average cooling rate of 5 ° C./s or more. This annealing is preferably continuous annealing performed in a continuous annealing line or a continuous hot dip galvanizing line in order to ensure cooling conditions and overaging treatment conditions.

  In annealing, it is necessary to heat to an annealing temperature of 800 ° C. or higher, which is in the α-γ2 phase region, for recrystallization and for the formation of the second phase. On the other hand, when the annealing temperature exceeds 950 ° C., the recrystallized grains are remarkably coarsened and the characteristics are remarkably deteriorated. The heating rate during heating need not be limited. For example, if the average heating rate from 300 to 700 ° C is less than 1 ° C / s, strain energy is released by recovery before recrystallization. Since the driving force for recrystallization tends to be reduced, it is preferably set to 1 ° C./s or more. In addition, the upper limit of the average temperature increase rate to 300-700 degreeC is about 50 degreeC / s in the present installation. Further, from 700 ° C. to the annealing temperature, it is preferable to heat at an average temperature increase rate of 0.1 ° C./s or more from the viewpoint of forming a recrystallized texture. On the other hand, when heating from 700 ° C to the annealing soaking temperature at an average heating rate of 20 ° C / s or more, transformation from the non-recrystallized part or non-recrystallized easily proceeds, and in terms of texture formation Since it tends to be disadvantageous, it is preferable to heat at an average heating rate of 20 ° C./s or less.

  The cold-rolled steel sheet after heating is separated into an α phase and a γ phase during cooling, and from the viewpoint of concentrating C in the γ phase, the temperature range from the annealing temperature to 580 ° C, which greatly affects this, is 5 to 5 ° C. It needs to be cooled at an average cooling rate of 30 ° C / s. If the average cooling rate is less than 5 ° C / s, the second phase is difficult to form, and a ferrite single phase structure is formed, and a TS of 440 MPa or more cannot be obtained. On the other hand, when the average cooling rate exceeds 30 ° C./s, the structure obtained by the ferrite transformation becomes a low-temperature transformation structure having a high dislocation density, and the ductility is lowered.

  Subsequently, it is necessary to perform an overaging treatment with a holding time of 30 to 600 seconds at an overaging temperature of 480 to 580 ° C. Fig. 4 shows the relationship between overaging temperature and tensile property values. Nmax is 5 or more at overaging temperatures of 480 to 580 ° C, high El and high U-El are obtained, and TS x El is 20000 MPa ·% or more. , TS x U-El is over 10000 MPa ·%. When the overaging temperature is less than 480 ° C., Nmax is less than 5, and it is not possible to obtain good TS × El and TS × U-El. On the other hand, when the over-aging temperature is 580 ° C or higher, the γ phase is unstable and a sufficient second phase cannot be obtained, Nmax becomes small, and good TS × El and TS × U-El. I can't. In addition, if the retention time of the overaging treatment is less than 30 s, the γ phase is unstable and a sufficient second phase cannot be obtained, and if it exceeds 600 s, excessive cementite is generated, so the retention time is set to 30 to 600 s. There is a need.

  The cold-rolled steel sheet after over-aging treatment suppresses the formation of excessive cementite phase, so that the average cooling rate is 5 ° C / s or higher to 100 ° C or lower, that is, at least 100 ° C, the average cooling rate is 5 ° C / s. Cooled as above.

  In addition, cooling after heating or cooling after overaging treatment can be performed by roll cooling, gas jet cooling, water quenching cooling, or the like.

  A plated layer may be formed on the cold-rolled steel sheet after annealing by electroplating or hot dipping. Further, the annealing may be performed in a continuous annealing line, once cooled to room temperature, and then hot dip galvanized in a hot dip galvanizing line, or further alloyed. Here, the plating layer is not limited to pure zinc and zinc-based alloy plating, but may of course be various plating layers conventionally applied to the steel sheet surface, such as Al or Al-based alloy plating.

  The cold-rolled steel plate or plated steel plate thus produced may be subjected to temper rolling or leveler processing for the purpose of shape correction and surface roughness adjustment. The total elongation of temper rolling or leveler processing is preferably in the range of 0.2 to 15%. This is because if it is less than 0.2%, the purpose of shape correction and surface roughness adjustment may not be achieved, and if it exceeds 15%, a significant reduction in ductility is caused. In addition, although the processing form differs between temper rolling and leveler processing, it has been confirmed that there is no significant difference between the two. In addition, temper rolling and leveler processing are effective even after plating.

  Steels A to Q having the compositions shown in Table 1 were melted in a converter and made into slabs by a continuous casting method. These slabs were heated to 1250 ° C., roughly rolled into sheet bars, and then hot-rolled steel sheets under the hot-rolling conditions shown in Tables 2 and 3. These hot-rolled steel sheets are pickled and cold-rolled at a reduction rate of 70% to form cold-rolled steel sheets (thickness: 1.2 mm), followed by annealing in the continuous annealing line under the annealing conditions shown in Tables 2 and 3. Then, temper rolling with an elongation of 0.5% was performed to prepare steel plate Nos. 1 to 50 samples. During annealing, the average temperature increase rate from 300 to 700 ° C. was 14 to 16 ° C./s, and the average temperature increase rate from 700 ° C. to the annealing temperature was 3 to 4 ° C./s. And about the obtained sample, the structure | tissue, a tensile characteristic, r value, and a texture were investigated by the following method.

  Microstructure: Observe the cross section of the plate parallel to the rolling direction using an optical microscope or scanning electron microscope. Using the above image analysis software, the area ratio of the ferrite phase, the area ratio of the martensite phase and the second phase, and the ferrite average The particle size and Nmax were determined.

  Tensile properties: JIS No. 5 tensile test specimens were taken from the sample at 90 ° to the rolling direction, and subjected to a tensile test at a crosshead speed of 10 mm / min in accordance with the provisions of JIS Z 2241. YS, TS, El, TS × El, U-El, and TS × U-El were obtained.

r value: JIS No. 5 tensile specimen taken from the sample in the rolling direction, 45 ° direction to the rolling direction, and 90 ° direction to the rolling direction, and the width of each specimen when 10% uniaxial tensile strain was applied. Strain and plate thickness strain were measured, and an average r value (average plastic strain ratio) was calculated from the following formula in accordance with the provisions of JIS S 2254 to evaluate deep drawability.
Average r value = (r ₀ + 2r ₄₅ + r ₉₀ ) / 4
Here, r _0, r _{45, and} r ₉₀ are plastic strain ratios measured with test pieces taken from 0 °, 45 °, and 90 ° directions with respect to the rolling direction, respectively.

Texture: Energy dispersive X-ray diffraction was performed at 1/4 thickness position of the sample. The measurement surface is the main diffraction surface of α-Fe (110), (200), (211), (220), (310), (222), (321), (400) Surface, (411) surface, (420) surface a total of 10 surfaces, relative intensity ratio with the non-directional standard sample to determine the X-ray diffraction integrated intensity ratio of each surface, (222) surface, ( The X-ray diffraction integrated intensity ratios P _(222), P _(200), P ₍₁₁₀₎ and P ₍₃₁₀₎ of the (200) plane, (110) plane and (310) plane were determined, and TA was calculated from the following equation.
TA = P ₍₂₂₂₎ / {P ₍₂₀₀₎ + P ₍₁₁₀₎ + P ₍₃₁₀₎ }
In addition, it can be said that the larger the TA, the more the texture that is more advantageous for the r value is developed.

  The results are shown in Tables 4, 5, and 6. In each of the present invention examples, Nmax is 5 or more, TS of 440 MPa or more, average r value of 1.2 or more, TS × El of 20000 MPa ·% or more, TS × U-El of 10,000 MPa ·% or more are obtained. I understand that.

FIG. 3 is a diagram showing an example of a SEM image (a) of a plate thickness section, a ferrite phase (b), a second phase (c), and a distribution (d) of the center of gravity of the second phase. It is a figure which shows the method of calculating | requiring the number of gravity centers of a 2nd phase. It is a figure which shows an example of the relationship between the number of gravity centers of a 2nd phase, and its frequency. It is a figure which shows the relationship between overaging treatment temperature and a tensile characteristic value.

Claims (4)

In mass%, C: 0.020 to 0.050%, Si: 1.0% or less, Mn: 1.0 to 2.5%, P: 0.005 to 0.1%, S: 0.01% or less, Al: 0.005 to 0.5%, N: 0.01% or less, In a matrix containing Nb: 0.010 to 0.3%, the contents of Nb and C satisfy the following formula (1), the balance is composed of Fe and inevitable impurities, and is composed of a ferrite phase, The second phase containing a martensite phase of 1% or more in terms of the area ratio relative to the entire structure has an area ratio of 15% or less, and the Nmax defined below is 5 or more and an average r value of 1.2 or more High strength steel plate with excellent ductility and deep drawability, characterized by
(Nb / 93) / (C / 12) = 0.2 to 0.7 (1)
In the formula (1), each element symbol represents the content (% by mass) of each element,
Nmax is a distribution diagram of the center of gravity of the second phase obtained by image processing the image observed by the scanning electron microscope of the steel plate thickness cross section, and draws a circle with the ferrite average grain diameter as the radius around an arbitrary center of gravity, This is the number of centroids that appears most frequently when the operation for obtaining the number of centroids included in a circle other than the center of gravity of the circle is performed on all the centroids recognized in the observed image.   2. The high-strength steel sheet excellent in ductility and deep drawability according to claim 1, further comprising Cr: 0.5% or less in mass% in addition to the above composition. In mass%, C: 0.020 to 0.050%, Si: 1.0% or less, Mn: 1.0 to 2.5%, P: 0.005 to 0.1%, S: 0.01% or less, Al: 0.005 to 0.5%, N: 0.01% or less, A steel slab having a composition containing Nb: 0.010 to 0.3%, the contents of Nb and C satisfying the following formula (1), the balance being Fe and inevitable impurities, and finishing rolling at 800 ° C. or higher Hot-rolled at a temperature, made into a rolled hot-rolled steel sheet at a winding temperature of 400 to 720 ° C., then cold-rolled into a cold-rolled steel sheet by cold rolling at a reduction rate of 40% or more, and the cold-rolled steel sheet is 800 to 950 ° C. Then, the temperature range from the annealing temperature to 580 ° C. is cooled at an average cooling rate of 5 to 30 ° C./s, and after holding at an overaging temperature of 480 to 580 ° C. for 30 to 600 s, A method for producing a high-strength steel sheet excellent in ductility and deep drawability, characterized by cooling at least to 100 ° C. at an average cooling rate of 5 ° C./s or higher;
(Nb / 93) / (C / 12) = 0.2 to 0.7 (1)
In the formula (1), each element symbol represents the content (% by mass) of each element.   4. The method for producing a high-strength steel sheet excellent in ductility and deep drawability according to claim 3, wherein the steel slab further contains, by mass%, Cr: 0.5% or less.

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